Gradient hard alloy and application of gradient hard alloy in preparation of bur

Abstract

According to the medium-high cobalt gradient hard alloy material provided by the invention, a gradient structure is formed in situ by calcining raw materials, the surface layer part of the gradient hard alloy material does not contain a brittle Ti cubic phase and a Zr-containing cubic phase, the gradient hard alloy material has high toughness, the subsurface layer part contains the Ti-containing cubic phase, the interior contains the Ti-containing cubic phase and the Zr-containing cubic phase, and the hardness and the plastic deformation resistance are high; and the plastic deformation requirement and the hardness requirement of the dental bur under the contact collision action with high-hardness enamel on the surface of irregular teeth in the use process can be met.

Description

TECHNICAL FIELD

[0001] The present application belongs to the technical field of hard alloy materials, and particularly relates to a gradient hard alloy, a preparation method thereof and application. BACKGROUND

[0002] A dental bur is a long rod-shaped rotary cutting instrument, and bur materials usually include steel, diamond and hard alloy. Among them, hard alloy materials are mainly formed by tungsten carbide powder and metal cobalt powder and other raw materials through powder metallurgy technology, and have the characteristics of high hardness, high compressive strength and high temperature resistance. In recent years, due to the development of precision processing industry, the working part of rotary cutting instruments is more and more manufactured by using hard alloy, and the bur is also developing from steel bur to hard alloy bur. The biggest feature of hard alloy bur is high hardness, and the hardness of hard alloy material is 2 times that of carbon steel, and is higher than that of teeth, which is 6 times that of tooth enamel. Hard alloy bur is more stable in cutting than diamond bur, produces less heat, and the working blade does not deform at high temperature, and has the characteristics of high durability. Hard alloy bur is different from diamond bur or steel bur, which does not grind the part to be removed of teeth during work, but cuts the tooth tissue to be removed, and the cut tooth surface is smooth and has no scratches. The working part of hard alloy bur can be used as reference for industrial rotary blades such as drill bits and end mills, and complex structures with complex shape, sharp blade and high cutting efficiency can be designed. Since the hard alloy bur is very sharp, it does not need to use gravity during use, and can work by relying on the weight of a dental high-speed handpiece, which is one of the advantages of hard alloy bur compared with steel bur or diamond bur.

[0003] The structure of the hard alloy turning tool is continuously improved and optimized, and the application is continuously developed. The patent for invention with the application publication number CN108577987 discloses a hard alloy turning tool for dental rotary instruments. It comprises a handle and a working part. The handle is connected with the dental rotary instrument, and the working part is connected with the handle. The working part is provided with an end blade and a chip removal groove. The end blade of the working part is fitted with the rake face cutting angle of the end blade by fitting the dentin stripping process with a double cubic B-spline curve. The invention can automatically chip by cutting towards the center of a parabola, can collect cutting fluid to the center point, and can accelerate the cooling and chip removal functions of cutting. By setting the blade angle, the built-up edge phenomenon is avoided on the rake face blade surface, and the tool life is prolonged. The patent for invention with the application publication number CN109551057A provides a turning tool for the final finishing processing of jewelry processing. The turning tool comprises a needle handle for installation and a machining head for processing. The needle handle is made of stainless steel material, and the machining head is made of hard alloy material. The needle handle and the machining head are welded and connected. The machining head comprises a cutting section and a transition section. The range of the conical angle of the cutting section is 4°-8°, the number of cutting edges is 6-10, and the spiral angle of the cutting edge is 10°-20°. The high hardness and wear resistance of the hard alloy material machining head is used to process jewelry, the machining head is more sharp, the service life is longer, and the surface of the jewelry is more smooth after processing.

[0004] The above invention is mainly aimed at the progress of the shape structure, cutting edge and angle, and application of the hard alloy turning tool, and there is no innovation and improvement of the hard alloy material used for the turning tool. The ordinary uniform hard alloy material is used. The turning tool usually uses YG hard alloy material (WC-Co). YG hard alloy has high transverse rupture strength and impact toughness. Especially when the cobalt content of the hard alloy is high, the strength and toughness are higher. High toughness can resist crack formation and expansion caused by impact during rotary cutting, prevent the cutting edge from appearing collapse, and prevent the turning tool from being damaged. Therefore, from the perspective of toughness, medium and high cobalt YG hard alloy is suitable as the material of the slender rod-shaped turning tool.

[0005] During tooth extraction, the patient's mouth is mostly open with all strength, in an uncomfortable state, so as to reduce the tooth extraction time of the patient, reduce the discomfort, the efficiency of the hard alloy dental bur resection tooth is required more and more. The minimum no-load speed of the gas turbine dental high-speed handpiece is 160,000 revolutions per minute, generally 160,000-600,000 revolutions per minute, such as the Austrian WH high-speed handpiece speed of 330,000 revolutions per minute, especially when removing impacted teeth, compared with ordinary treatment, the cutting efficiency is required to be high, and the speed is required to be high. The high speed makes the bur working end and the tooth rub quickly to generate a large amount of heat energy, and there are many tooth fragments during work, under the condition of such high speed and fragment blockage, not only the toughness of the hard alloy material is required to be high, but also the hard alloy material must be required to have high hardness and strong plastic deformation resistance. However, the medium-high cobalt YG hard alloy has low hardness and insufficient plastic deformation resistance, which hinders the cutting rate upgrading of the hard alloy bur. In addition, the gradient hard alloy raw material contains cubic phase compounds such as TiN and ZrC. Since the cubic phase compound belongs to sodium chloride type face-centered cubic (f.c.c) structure, the interatomic gap of the face-centered cubic lattice is large, and the oxygen solubility is strong. The cubic phase carbide is easy to oxidize during preparation. The oxidation reaction of the cubic phase compound increases the metal oxide in the raw material and increases the combined oxygen content. Oxidation causes the wettability of the hard phase and the binder to decrease during sintering, thereby hindering sintering densification, resulting in a series of problems such as increased porosity, reduced strength and hardness of the gradient hard alloy. SUMMARY

[0006] To solve the above problems, the present application provides a medium-high cobalt gradient hard alloy material, which uses raw materials Co powder, TiN powder, ZrC powder, W powder and WC powder, calcines, and forms a gradient structure in situ. The surface layer of the gradient hard alloy material only contains metal elements W and Co, does not contain brittle Ti cubic phase and Zr cubic phase, has high toughness, the subsurface layer contains Ti cubic phase, the inside contains Ti cubic phase and Zr cubic phase, the hardness and plastic deformation resistance of the subsurface layer and the inside are high, and can meet the plastic deformation requirement and hardness requirement of the dental bur under the action of the collision of the irregular tooth surface with the high-hardness tooth enamel during use.

[0007] The present application aims to provide a gradient hard alloy, which comprises a surface layer, a subsurface layer and an inside from outside to inside, the metal elements of the surface layer comprise W and Co elements, the metal elements of the subsurface layer comprise W, Co and Ti elements, and the metal elements of the inside comprise W, Co, Ti and Zr elements,

[0008] In the gradient hard alloy, the content of Co element is greater than 10wt%,

[0009] The surface layer portion has a thickness of 60-110 μm.

[0010] The present application also aims to provide a gradient cemented carbide, which is prepared by grinding Co powder, cubic phase powder, W powder and WC powder to form a green body, and then calcining the green body under vacuum and in a protective atmosphere.

[0011] The present application also aims to provide a method for preparing a gradient cemented carbide, which is prepared by grinding Co powder, cubic phase powder, W powder and WC powder to form a green body, and then calcining the green body under vacuum and in a protective atmosphere.

[0012] The present application also aims to provide a use of the gradient cemented carbide for preparing a dental tool bur.

[0013] The present application has the following advantages:

[0014] (1) Cubic phase compounds such as TiN and ZrC are usually used in the raw materials of cemented carbide. Cubic phase compounds belong to sodium chloride type face-centered cubic (f.c.c) structure, and the interstitial space of the face-centered cubic lattice is large and has strong oxygen solubility. Cubic carbide is easy to oxidize during preparation, and the oxidation reaction of cubic phase compounds increases the content of metal oxides in the raw materials and increases the combined oxygen content. Oxidation causes the wettability of the hard phase and the binder to decrease during sintering, thereby hindering sintering densification, resulting in a series of problems such as increased porosity, reduced strength and hardness of the gradient cemented carbide. The present application reduces the oxygen content of the cubic phase compound raw material powder through atmosphere protection heat treatment, thereby reducing the porosity and improving the strength and hardness.

[0015] (2) The surface layer of the gradient cemented carbide material of the present application has high toughness, and the internal performance has strong resistance to plastic deformation, so the material as a whole is resistant to impact vibration and plastic deformation.

[0016] (3) The gradient cemented carbide material of the present application does not have a sudden interface between the surface layer and the interior, avoiding the problems of insufficient bonding strength of the sudden interface and interface stress caused by the difference in physical properties such as thermal expansion coefficient and elastic modulus on both sides of the interface.

[0017] (4) The gradient structure of the gradient cemented carbide material of the present application is formed in situ during sintering of the alloy, does not need to be layered, does not increase additional processes, is simple in process, easy to operate and implement, and has readily available raw materials and low cost.

[0018] (5) The present application reduces the oxygen content of the cubic phase powder through atmosphere protection heat treatment, and the deoxidation effect is significantly better than that of the traditional powder metallurgy method for preparing gradient cemented carbide. BRIEF DESCRIPTION OF DRAWINGS

[0019] Figure 1 SEM image showing the microstructure of the near-surface region of the gradient cemented carbide I in Example 1;

[0020] Figure 2 EDS map showing the Ti element surface distribution of the gradient cemented carbide I in Example 1;

[0021] Figure 3 EDS map showing the Zr element surface distribution of the gradient cemented carbide I in Example 1;

[0022] Figure 4 Composition element map showing the surface region of the gradient cemented carbide I in Example 1;

[0023] Figure 5 Composition element map showing the subsurface region of the gradient cemented carbide I in Example 1;

[0024] Figure 6 Composition element map showing the internal region of the gradient cemented carbide I in Example 1;

[0025] Figure 7 High magnification SEM image showing the internal region microstructure of the gradient cemented carbide I in Example 1;

[0026] Figure 8 Composition element map showing the core phase in the internal region of the gradient cemented carbide I in Example 1;

[0027] Figure 9 Composition element map showing the ring phase in the internal region of the gradient cemented carbide I in Example 1;

[0028] Figure 10 Metallographic structure map showing the gradient cemented carbide I in Example 1;

[0029] Figure 11 SEM image showing the microstructure of the near-surface region of the gradient cemented carbide II in Example 2;

[0030] Figure 12 SEM image showing the microstructure of the near-surface region of the cemented carbide A in Comparative Example 1;

[0031] Figure 13 SEM image showing the microstructure of the near-surface region of the cemented carbide B in Comparative Example 2;

[0032] Figure 14 SEM image showing the microstructure of the near-surface region of the cemented carbide C in Comparative Example 3;

[0033] Figure 15 EDS map showing the Ti element surface distribution of the cemented carbide C in Comparative Example 3;

[0034] Figure 16EDS map of Zr element surface distribution of hard alloy C in Comparative Example 3 is shown.

[0035] Figure 17 A metallographic structure diagram of hard alloy D in Comparative Example 4 is shown. DETAILED DESCRIPTION

[0036] The present application will be described in detail below through specific embodiments, and the features and advantages of the present application will become more apparent and explicit with these descriptions.

[0037] The present application provides a gradient hard alloy, which comprises a surface layer part, a subsurface layer part and an inner part from outside to inside, the metal elements of the surface layer part comprising W and Co elements, the metal elements of the subsurface layer part comprising W, Co and Ti elements, and the metal elements of the inner part comprising W, Co, Ti and Zr elements.

[0038] In the gradient hard alloy, the content of Co element is greater than 10wt%, preferably 13-20wt%, and more preferably 13-18wt%. The turning needle hardness of the existing medium-high cobalt YG hard alloy is low, and the plastic deformation resistance is not enough, which hinders the cutting speed upgrade of the hard alloy turning needle. In the present application, appropriately increasing the content of cobalt in the gradient hard alloy is beneficial to improve the toughness and processability of the material, and more importantly, in cooperation with other raw materials and raw material particle size, the thickness of the surface layer part can be increased, the machining allowance can be increased, the turning needle hardness and plastic deformation resistance can be improved, and the cutting speed of the turning needle can be improved.

[0039] In the surface layer part, Ti element and Zr element are not contained, and due to the limitations of detection conditions and error background, the detected content of Ti element is less than 0.001wt%, and the detected content of Zr element is less than 0.001wt%. In the subsurface layer part, Zr element is not contained, and due to the limitations of detection conditions and error background, the detected content of Zr element is less than 0.001wt%.

[0040] The thickness of the surface layer part is 60-110μm, preferably 65-100μm, and more preferably 70-90μm. Since the surface layer is a high-toughness WC-Co composition, the increase of the thickness of the surface layer of the gradient alloy means that the proportion of the high-toughness part is increased, and the overall toughness of the alloy is increased. Since the teeth are irregular in shape and uneven, the turning needle is not continuously cut when cutting the teeth, but is intermittently cut and frequently subjected to impact force. The high toughness of the turning needle material can resist impact, so that the turning needle will not collapse, thereby improving the service life of the turning needle. At the same time, the increase of the thickness of the surface layer part increases the machining allowance range when machining the turning needle, which is convenient for machining and application.

[0041] The total thickness of the surface layer and the subsurface layer is 80-120 microns, preferably 80-110 microns, and more preferably 85-100 microns. There is only Ti-containing cubic phase in the subsurface layer, and there is no Zr-containing cubic phase. Compared with the inside, the toughness of the subsurface layer is higher than that of the inside. Therefore, the increase in the thickness of the subsurface layer means that the proportion of the part with higher toughness increases, and the overall toughness of the alloy increases. Since the teeth are irregular in shape and uneven, the cutting needle is not continuously cut when cutting the teeth, but is intermittently cut and frequently subjected to impact force. The cutting needle with high toughness can resist impact, so that the cutting needle will not collapse, thereby improving the service life of the cutting needle. The surface layer does not contain Ti cubic phase and Zr cubic phase, the subsurface layer contains Ti cubic phase, and the inside contains Ti cubic phase and Zr cubic phase. The gradient hard alloy of the application only contains metal elements W and Co in the surface layer, and does not contain brittle Ti cubic phase and Zr cubic phase, so it has high toughness. The subsurface layer has Ti cubic phase, and the inside has Ti cubic phase and Zr cubic phase. Since the cubic phase has high hardness, the hardness and plastic deformation resistance of the subsurface layer and the inside are high.

[0042] The inside of the gradient hard alloy material has a core phase and a ring phase, and the ring phase wraps the core phase. The Zr content in the core phase is higher than that in the ring phase, and the Ti content in the ring phase is higher than that in the core phase.

[0043] The density of the gradient hard alloy is greater than 12 g / cm 3 , the transverse rupture strength is greater than 2200 MPa, the hardness (Hv) is 1200 MPa, and the relative magnetic saturation strength is greater than 18 Gs·cm 3 / g.

[0044] Preferably, the porosity of the gradient hard alloy is A02B02C00. In the field of hard alloys, pores are divided into three types: A, B and C. Type A pores are less than 10 microns, and are divided into A02, A04, A06 and A08 levels, which respectively represent a pore volume ratio of 0.02%, 0.06%, 0.2% and 0.6%. Type B pores are 10-25 microns, and are represented by B02, B04, B06 and B08, which respectively represent the proportion of pores of this size as 140 / cm 2 , 430 / cm 2 , 1300 / cm 2 , 4000 / cm 2 . Type C pores refer to graphite inclusions, which are divided into C02, C04, C06 and C08 levels.

[0045] The application provides a gradient hard alloy. Co powder, cubic phase powder, W powder and WC powder are ground to form a green body, which is calcined under vacuum and protective atmosphere conditions in sequence.

[0046] The cubic phase powder is obtained by mixing and grinding TiN powder (titanium nitride powder) and ZrC powder (zirconium carbide powder), pre-sintering under vacuum, and then calcining under a protective atmosphere. The weight of the cubic phase powder is 2-10 parts, preferably 3-8 parts, and more preferably 4-6 parts. The cubic phase in the present application includes ZrC and TiN. When the addition amount of ZrC and the component ratio of Ti in the raw material reach a certain range, the hard phase will be dispersedly distributed in the matrix, playing a strong solid solution strengthening role and improving the hardness and strength of the material. The main role of TiN is to act as a nitrogen source. Through decomposition at high temperature, the thermodynamic coupling of N and the cubic phase metal elements promotes element diffusion, thereby forming a gradient structure. When the content of the cubic phase is very low, the effect is not obvious. When the content of the cubic phase is too high, the cubic phase is a brittle phase, which will lead to the decrease of the strength and toughness of the whole cemented carbide. Therefore, the content needs to be controlled within a certain range.

[0047] The present application also provides a preparation method of the gradient cemented carbide. The method grinds Co powder, cubic phase powder, W powder (tungsten powder), and WC powder (tungsten carbide powder) to form a green body, and then calcines the green body under vacuum and a protective atmosphere to obtain the gradient cemented carbide.

[0048] The cubic phase powder is obtained by mixing and grinding TiN powder (titanium nitride powder) and ZrC powder (zirconium carbide powder), pre-sintering under vacuum, and then calcining under a protective atmosphere.

[0049] The average particle size of the TiN powder is 1.5-4.0 μm, preferably 2.0-3.5 μm, and more preferably 2.5-3.0 μm, and the weight fraction is 1-5, preferably 1.5-4, and more preferably 2-3. In the prior art, TiN is directly added, and the nitrogen activity in the matrix is high. If one-step sintering process is used, TiN tends to decompose and generate nitrogen gas before the gradient sintering temperature is reached, thereby reducing the density of the alloy. Therefore, in order to avoid premature decomposition of the nitrogen-containing phase, it is generally necessary to first pre-sinter at a lower temperature, and introduce nitrogen protection during the sintering process. After cooling to obtain a homogeneous structure of the cemented carbide, gradient sintering is performed. In the present application, the sintering process is carried out under vacuum without introducing nitrogen protection. The smaller the particle size of the TiN powder, the greater the sintering activity, and the more unstable and easier to decompose TiN is. In the present application, the particle size of the TiN raw material is not ultra-fine or sub-micron, but medium, so as to avoid premature decomposition of TiN. In the case where ZrC is added, within a certain range, the thickness of the gradient surface layer increases with the increase of the TiN content. In the case where ZrC is not added, the thickness of the gradient surface layer decreases with the increase of the TiN content. The smaller the particle size of TiN, the greater the sintering activity, and the faster the diffusion under the same sintering temperature and sintering time, thereby increasing the thickness of the gradient surface layer. However, the prerequisite is that TiN cannot be prematurely decomposed before the sintering temperature.

[0050] The average particle size of the ZrC powder is 0.3-1.2 μm, preferably 0.4-1.0 μm, and more preferably 0.5-0.8 μm, and the weight fraction is 1-6, preferably 1.5-5, and more preferably 2-4. ZrC will increase the liquidus temperature of the system, inhibit the diffusion of W and other metal elements, thereby reducing the amount of intermetallic phase, refining the structure, and strengthening the matrix material. However, when the ZrC content exceeds a certain value, ZrC will affect the wettability of the binder phase to the ceramic phase, weaken the binding effect, and cause micro-pores and micro-cracks in the material matrix, thereby reducing the density of the material and adversely affecting the hardness and transverse rupture strength of the material. The finer the particle size of ZrC, the more significant the promotion effect on the hardness of the alloy, and the finer the particle size, the more conducive to the diffusion of Zr element, and the slightly increased thickness of the gradient surface layer.

[0051] The TiN powder and the ZrC powder are mixed and ground by using Φ6 mm WC-6wt%Co cemented carbide balls, and the ball-to-material weight ratio is 10:(1-1.5), and the grinding speed is 260-360 r / min for 1-2 h.

[0052] The pre-burning under vacuum is to mix the ground TiN powder and ZrC powder mixture at 90-130℃, vacuum degree of 0.5-1.5Pa, and heat preservation for 1-4h, preferably at 100-120℃, vacuum degree of 0.5-1Pa, and heat preservation for 1-2h.

[0053] The calcining under protective atmosphere is to fill in the mixed gas of nitrogen and argon at 1100-1400℃, pressure of the mixed gas is 1-8MPa, pressure ratio of nitrogen to argon is 1:10-30, heat preservation time is 1-4h, and then cooling in the furnace to obtain cubic phase powder, preferably filling in the mixed gas of nitrogen and argon at 1200-1300℃, pressure of the mixed gas is 1-5MPa, pressure ratio of nitrogen to argon is 1:15-25, heat preservation time is 1-2h, and then cooling in the furnace to obtain cubic phase powder.

[0054] The pre-burning and calcining under protective atmosphere can reduce the oxygen content of cubic phase raw material, thereby reducing the porosity in the subsequent sintering process.

[0055] The average particle size of Co powder is 0.6-2.1μm, preferably 0.8-1.7μm, more preferably 1-1.3μm, and the weight fraction is greater than 10, preferably 13-20, more preferably 13-18. Increasing the Co content increases the element diffusion channel, so under the same sintering conditions, the gradient surface layer becomes thicker. Since the metal Co is more flexible and has lower hardness than the ceramic hard phase, increasing the Co content also improves the flexibility of the alloy as a whole, but the hardness of the alloy decreases.

[0056] The selection of raw material particle size has an effect on the performance of the hard alloy. The particle size of the raw material refers to the granularity, and a raw material particle contains a large number of grains. The grain size, which is essentially the grain size, has an effect on the performance of the hard alloy. The finer the grain size, the higher the hardness of the hard alloy. Generally speaking, fine-grained raw materials contain fine grains, and coarse-grained raw materials contain coarse grains. After grinding, the particle size of the particles decreases, which only means that the number of grains contained in the particles decreases, and the size of the grains contained in the particles does not decrease significantly. Therefore, the selection of raw material particles means the control of grain size.

[0057] The average particle size of the W powder is 1-6 μm, preferably 1.5-5 μm, and more preferably 2-4 μm, and the weight fraction is 0.1-1.0, preferably 0.2-0.8, and more preferably 0.4-0.6. When rubber is used as the forming agent for the hard alloy, due to the degradation of the polymer constituting the rubber during the removal of the forming agent, partial cracking of the polymer inevitably occurs, and the cracking products include a certain amount of carbon black, which remains in the alloy and causes carbonization. In order to reduce the carbonization caused by the cracking of the forming agent and the resulting change in the composition, the present application introduces elemental W powder to balance the excess carbon content caused by the cracking of the forming agent. The combination of W and C forms WC, and a deficiency of W results in a relatively large amount of residual C, which causes the hard alloy to be carburized and the performance to be poor, while a high content of W results in a relatively low proportion of C, which causes the hard alloy to be decarburized, brittle, and poor in performance.

[0058] The average particle size of the WC powder is 3-9 μm, preferably 4-8 μm, and more preferably 5-7 μm, and the weight fraction is 70-95, preferably 75-90, and more preferably 80-85. WC is the main hard phase of the hard alloy and is the basic component of the hard alloy. When the content of WC is higher and the relative content of Co is lower, the hardness of the hard alloy increases, and the toughness decreases. The finer the particle size of WC, the higher the hardness of the hard alloy, the higher the strength, and the thicker the gradient surface layer.

[0059] The grinding is performed on the mixture of Co powder, TiN powder, ZrC powder, W powder, and WC powder using a hard alloy ball, and the weight ratio of the hard alloy ball to the mixture is 5-15:1, and preferably 8-12:1. Preferably, the hard alloy ball is a Φ8 mm WC-6wt%Co sub-micro hard alloy ball. The grinding speed is 50-90 r / min, and the grinding time is 24-96 h.

[0060] The grinding medium is preferably anhydrous ethanol. After the grinding is completed, the mixed slurry is filtered through a 300-500 mesh screen. Drying is performed, for example, under a temperature of 85-120 °C and a vacuum degree of 1-5 Pa. A rubber forming agent, such as an SD rubber forming agent, is added, and the green body is pressed, preferably under a pressure of 300-500 MPa. The amount of the rubber forming agent added is 4-7 parts by weight.

[0061] The calcination temperature under vacuum is 450-1450 °C, preferably 500-1400 °C, and more preferably 550-1350 °C, and the calcination time is 1.5-4.5 h, preferably 2-4 h, and more preferably 2.5-3.5 h.

[0062] Preferably, the forming agent is removed at a temperature of 400-700°C, preferably 500-600°C, at a heating rate of 0.5-2.5°C / min, for 0.5-3h, preferably 1h, under a vacuum of 10-15Pa, preferably 12-18Pa; then the solid phase sintering is completed at a sintering temperature of 1150-1250°C, preferably 1200-1250°C, for 0.5-2h, such as 1h, at a heating rate of 2-6°C / min, under a vacuum of 5-10Pa, preferably 8-12Pa; and then the liquid phase sintering is completed at a sintering temperature of 1280-1350°C, preferably 1300-1350°C, for 30-90min, preferably 25-40min, at a heating rate of 1-5°C / min, preferably 2-4°C / min, under a vacuum of 1-5Pa, such as 4-5Pa.

[0063] The calcination temperature under the protective atmosphere is 1300-1550°C, preferably 1350-1500°C, and the calcination time is 30-90min, preferably 40-80min, more preferably 50-70min. The protective atmosphere is preferably argon, and the argon pressure is 5-10MPa, preferably 5-6MPa. The heating rate is 2-8°C / min, preferably 3-5°C / min.

[0064] During the high-temperature sintering process, the TiN in the surface layer of the hard alloy decomposes in a denitrogenation environment, such as a protective atmosphere, and releases N2 into the surrounding environment. Due to the thermodynamic coupling between elements Ti and N and between Zr and N, the reduction of the N content in the surface layer causes the elements Ti and Zr in the surface layer to migrate towards the interior to regions with high N activity, resulting in the removal of the brittle cubic phase TiN and ZrC in the surface layer and the migration of the element Ti in the surface layer to the subsurface layer. Since the migration and diffusion ability of Zr is greater than that of Ti, the Zr elements in the surface layer and the subsurface layer both migrate to the interior, and the region where Zr is removed is farther than the region where Ti is removed. The raw materials for preparing the hard alloy are uniformly mixed, and the hard alloy forms a gradient structure in situ during the sintering process.

[0065] The cubic phase compounds, such as TiN and ZrC, contained in the raw materials are prone to oxidation during the preparation process because cubic phase compounds belong to sodium chloride type face-centered cubic (f.c.c) structure, the interatomic spacing of the face-centered cubic lattice is large, and the oxygen solubility is strong. The oxidation reaction of the cubic phase compounds leads to an increase in the metal oxides and the combined oxygen content in the raw materials. Oxidation causes the wettability of the hard phase and the binder to decrease during sintering, thereby hindering sintering densification, resulting in a series of problems such as an increase in the porosity of the gradient hard alloy and a decrease in the strength and hardness. The present application reduces the oxygen content of the cubic phase compound raw material powder through atmosphere protection heat treatment, thereby reducing the porosity, improving the strength and hardness, and meeting the high requirements of dental burs on the comprehensive performance of hard alloy materials.

[0066] The present application also provides the use of the gradient cemented carbide for preparing dental tools.

[0067] The gradient cemented carbide material of the present application has high toughness of the surface layer, and the high toughness surface layer can resist crack formation and propagation caused by impact, thus being suitable for dental burs used under high speed rotation, impact and vibration conditions. Meanwhile, the gradient cemented carbide has high hardness and strong resistance to plastic deformation in the sub-surface layer and the inner part, thus being able to resist plastic deformation under the impact of high hardness enamel of irregular tooth surface during use of the dental bur.

[0068] Examples

[0069] Example 1

[0070] TiN powder and ZrC powder were mixed in a stirring ball mill at a mass ratio of 1:1, the average particle size of the TiN powder was 2.7 μm, the average particle size of the ZrC powder was 0.6 μm, the grinding ball was a Φ6 mm WC-6wt%Co cemented carbide ball, the ball-to-material weight ratio was 10:1, and the grinding was carried out at a speed of 300 r / min for 1 h. The mixed TiN powder and ZrC powder after grinding were put into a graphite crucible, and then sintered in an atmosphere sintering furnace at a temperature of 1100 °C and a vacuum degree of 0.5 Pa for 1 h; then a mixed gas of nitrogen and argon was filled in at a temperature of 1210 °C, the pressure of the mixed gas was 2 MPa, the pressure ratio of nitrogen to argon was 1:19, and the holding time was 1 h, and then the furnace was cooled down to obtain cubic phase powder. After the atmosphere protection heat treatment, the oxygen content was reduced from 1.44wt% to 0.25wt%.

[0071] The raw materials were weighed, the weight of the Co powder with an average particle size of 1.1 μm was 13, the weight of the cubic phase powder was 4, the weight of the W powder with an average particle size of 2.0 μm was 0.4, and the weight of the WC powder with an average particle size of 6.0 μm was 82.6.

[0072] Then, the mixed raw material powder was put into a drum-type ball mill for grinding, the grinding ball was a Φ8 mm WC-6wt%Co sub-micro cemented carbide ball, the ball-to-material weight ratio was 10:1, the grinding medium was anhydrous ethanol, the amount of which was 200 mL, and the grinding was carried out at a speed of 60 r / min for 48 h. After the grinding, the cemented carbide slurry was filtered through a 400 mesh screen, and then vacuum dried at 5 Pa and 90 °C. After drying, the SD rubber forming agent was added in a weight of 5.5, and then vacuum dried again at 5 Pa and 90 °C. The mixed material after drying was filtered through an 80 mesh screen, and then pressed into a green body under a pressure of 400 MPa.

[0073] The green blank was placed in a vacuum furnace. (1) The heating rate was 1.3℃ / min, and the temperature was held at 560℃ for 1h. The forming agent was removed under a vacuum of 15Pa. (2) The heating rate was 3.6℃ / min, and the temperature was held at 1210℃ for 1h. The solid-phase sintering was completed under a vacuum of 10Pa. (3) The heating rate was 2.5℃ / min, and the temperature was held at 1310℃ for 35min. The liquid-phase sintering was completed under a vacuum of 5Pa. (4) The vacuum-sintered cemented carbide was placed in a pressure sintering furnace. The pressure sintering heating rate was 4.2℃ / min, and the temperature was held at 1440℃ for 60min. The argon pressure was 5MPa. Gradient cemented carbide I was obtained.

[0074] Scanning electron microscope (SEM) images of the microstructure of graded cemented carbide I () Figure 1 ).Depend on Figure 1 It can be seen that the surface layer thickness of gradient cemented carbide I is about 71 μm, and the total thickness of the surface and subsurface layers is about 97 μm.

[0075] The EDS plot of Ti element surface distribution in gradient cemented carbide I shows that the Ti element content in the surface region is almost zero, while the Ti element density in the subsurface region is higher than that in the interior region. Figure 2 As shown.

[0076] The EDS plot of Zr surface distribution shows that the Zr content in the surface and subsurface regions is almost zero. Figure 3 As shown.

[0077] Elemental composition map of the surface region obtained from the test ( Figure 4 As can be seen, the only metallic elements in the surface layer are W and Co; the elemental composition of the subsurface region obtained from the test ( Figure 5 As can be seen from the subsurface metallic elements, there are W, Co, and Ti elements; the elemental composition of the internal region obtained from the test ( Figure 6 As can be seen, the internal metallic elements are W, Co, Ti, and Zr.

[0078] Figure 7 The image shows a high-magnification SEM image of the internal microstructure. In the internal microstructure, the white phase is the WC phase, the black phase is the Co-based bonding phase, and there are also dark gray core phases and light gray ring phases. Figure 8 The image shows the elemental composition of the core phase, which has a higher Zr content. Figure 9 This is an elemental diagram of the cyclic phase composition, showing that the cyclic phase has a higher Ti content.

[0079] The porosity of gradient cemented carbide I was determined to be A02B02C00, and its metallographic structure is shown in the figure below. Figure 10 As shown, no obvious pores are visible on its surface; its density is 13.04 g / cm³. 3, the transverse rupture strength is 2292 MPa, the hardness (Hv) is 1303 MPa, and the relative magnetic saturation strength is 18.93 Gs·cm 3 / g.

[0080] Example 2

[0081] The gradient cemented carbide II is prepared according to the method of Example 1, except that the mass ratio of TiN powder to ZrC powder is 1:1.5 when preparing the cubic phase powder, and the oxygen content is reduced from 1.68 wt% to 0.34 wt% after the atmosphere protection heat treatment. The weight fraction of Co powder with an average particle size of 1.1 μm is 13, the weight fraction of the cubic phase powder is 5, the weight fraction of W powder with an average particle size of 2.0 μm is 0.4, and the weight fraction of WC powder with an average particle size of 6.0 μm is 81.6.

[0082] The surface layer of the prepared gradient cemented carbide II has a thickness of about 67 μm, the total thickness of the surface layer and the subsurface layer is about 92 μm, and the SEM image of the microstructure is shown in Figure 11 .

[0083] The porosity of the gradient cemented carbide II is A02B02C00, the density is 12.81 g / cm 3 , the transverse rupture strength is 2395 MPa, the hardness (Hv) is 1282 MPa, and the relative magnetic saturation strength is 19.07 Gs·cm 3 / g.

[0084] Comparative Example

[0085] Comparative Example 1

[0086] The cemented carbide A is prepared according to the method of Example 1, except that no cubic phase powder is added, the weight fraction of Co powder with an average particle size of 1.1 μm is 13, the weight fraction of W powder with an average particle size of 2.0 μm is 0.4, and the weight fraction of WC powder with an average particle size of 6.0 μm is 86.6. The SEM image of the microstructure of the prepared cemented carbide is shown in Figure 12 . The porosity is A02B00C00, the density is 14.19 g / cm 3 , the transverse rupture strength is 2239 MPa, the hardness (Hv) is 1220 MPa, and the relative magnetic saturation strength is 18.69 Gs·cm 3 / g.

[0087] Compared with Example 1, the surface layer and the internal microstructure of the cemented carbide are consistent, and no composition gradient structure is formed, which indicates that the introduction of cubic phase TiN and ZrC is a necessary condition for the formation of the gradient structure of the cemented carbide in the present application.

[0088] Comparative Example 2

[0089] Hard alloy B was prepared according to the method of Example 1, except that when preparing the cubic phase powder, TiN was not added; the weight fraction of Co powder with an average particle size of 1.1 μm was 13, the weight fraction of cubic phase powder was 2, the weight fraction of W powder with an average particle size of 2.0 μm was 0.4, and the weight fraction of WC powder with an average particle size of 6.0 μm was 84.6.

[0090] The SEM image of the microstructure of the prepared hard alloy B is shown in Figure 13 The porosity of the alloy was A02B02C00, the density was 13.72 g / cm 3 , the transverse rupture strength was 2007 MPa, the hardness (Hv) was 1188 MPa, and the relative magnetic saturation was 16.96 Gs-cm 3 / g.

[0091] Compared with Example 1, the prepared hard alloy B did not form a gradient structure, indicating that the introduction of only cubic phase ZrC cannot form a gradient structure.

[0092] Comparative Example 3

[0093] Hard alloy C was prepared according to the method of Example 1, except that when preparing the cubic phase powder, the mass ratio of TiN powder to ZrC powder was 1:4; the weight fraction of Co powder with an average particle size of 1.1 μm was 13, the weight fraction of cubic phase powder was 10, the weight fraction of W powder with an average particle size of 2.0 μm was 0.4, and the weight fraction of WC powder with an average particle size of 6.0 μm was 76.6.

[0094] The surface layer thickness of hard alloy C was about 8.5 μm, the microstructure is shown in Figure 14 , the Ti element distribution is shown in Figure 15 , and the Zr element distribution is shown in Figure 16 .

[0095] The porosity of hard alloy C was A02B02C00, the density was 11.93 g / cm 3 , the transverse rupture strength was 2694 MPa, the hardness (Hv) was 891 MPa, and the relative magnetic saturation was 19.17 Gs-cm 3 / g.

[0096] Compared with Example 1, the surface layer thickness of the hard alloy was very thin, there was no obvious subsurface layer, and the hardness of the alloy was significantly reduced, indicating that when the ZrC addition amount is too high, the gradient surface layer thickness is thinned and is not conducive to the hardness of the alloy.

[0097] Comparative Example 4

[0098] The raw materials were weighed and mixed uniformly, wherein the Co powder with an average particle size of 1.1 μm was 13 parts by weight, the TiN powder with an average particle size of 2.7 μm was 2 parts by weight, the ZrC powder with an average particle size of 0.6 μm was 2 parts by weight, the W powder with an average particle size of 2.0 μm was 0.4 parts by weight, and the WC powder with an average particle size of 6.0 μm was 82.6 parts by weight.

[0099] Then, the mixed raw material powder was added to a drum-type ball mill for grinding, the grinding ball was a Φ8 mm WC-6wt%Co sub-micro cemented carbide ball, the ball-to-material weight ratio was 10:1, the grinding medium was anhydrous ethanol, the amount of which was 200 mL, and the grinding was carried out at a speed of 60 r / min for 48 h. After grinding, the cemented carbide slurry was filtered through a 400-mesh screen and vacuum dried at 5 Pa and 90°C. After drying, 5.5 parts by weight of SD rubber forming agent was added, and after mixing uniformly, the mixture was vacuum dried again at 5 Pa and 90°C. The dried mixture was filtered through an 80-mesh screen and pressed into a green body under a pressure of 400 MPa.

[0100] The green body was placed in a vacuum furnace, (1) the heating rate was 1.3°C / min, the temperature was kept at 560°C for 1 h, and the vacuum degree was 15 Pa to remove the forming agent; (2) the heating rate was 3.6°C / min, the temperature was kept at the sintering temperature of 1210°C for 1 h, and the vacuum degree was 10 Pa to complete the solid-phase stage sintering; (3) the heating rate was 2.5°C / min, the temperature was kept at the sintering temperature of 1310°C for 35 min, and the vacuum degree was 5 Pa to complete the liquid-phase stage sintering. (4) The vacuum-sintered cemented carbide was placed in a pressure sintering furnace, the pressure sintering heating rate was 4.2°C / min, the temperature was kept at 1440°C for 60 min, the argon pressure was 5 MPa, and the gradient sintering was completed. Cemented carbide D was obtained.

[0101] The prepared cemented carbide D had a metallographic structure as shown in Figure 17 The alloy porosity was A04B04C00, and obvious pores could be seen on the surface; the density was 12.78 g / cm 3 , the transverse rupture strength was 1978 MPa, and the hardness (Hv) was 1082 MPa.

[0102] Compared with Example 1, the porosity of Cemented Carbide D is increased, the density is decreased, the transverse rupture strength and the hardness are all decreased. The increase of porosity will inevitably decrease the density of the cemented carbide. The transverse rupture strength reflects the ability of the cemented carbide to resist deformation or fracture under the action of bending external force. Since the porosity is a kind of defect that can become a fracture source, the increase of porosity leads to the increase of the number of fracture sources, so the transverse rupture strength is decreased. The hardness value reflects the ability of the cemented carbide to resist local plastic deformation. Since the porosity has no supporting capacity for the material, the increase of the porosity makes the local area material under the pressure head during the hardness measurement tend to deviate to the porosity, so that the surface area of the indentation is increased, and the hardness is decreased. When the cubic phase raw material powder is not subjected to atmosphere protection heat treatment, since the cubic phase TiN, ZrC and the like belong to sodium chloride type face-centered cubic structure, the interatomic gap is large, and the oxygen dissolving capacity is strong, and it is also easy to be oxidized during preparation. If more oxygen cannot be removed in time at the low temperature stage of the gradient cemented carbide sintering, the oxygen will react with C in the gradient cemented carbide at the high temperature sintering stage to generate CO, CO2 and the like, leading to the increase of the porosity of the gradient cemented carbide, thereby leading to the decrease of the mechanical properties such as strength and hardness of the gradient cemented carbide. It is indicated that in the present system, the atmosphere protection heat treatment of the cubic phase raw material powder can reduce the oxygen content in the powder, thereby improving the mechanical properties of the gradient cemented carbide.

[0103] Comparative Example 5

[0104] The preparation process of Example 1 is repeated, except that the cubic phase raw material powder TiN is not subjected to atmosphere protection heat treatment.

[0105] The prepared gradient cemented carbide has a porosity of A04B02C00, a density of 12.91 g / cm 3 , a transverse rupture strength of 2057 MPa, and a hardness (Hv) of 1205 MPa.

[0106] Comparative Example 6

[0107] The preparation process of Example 1 is repeated, except that the cubic phase raw material powder ZrC is not subjected to atmosphere protection heat treatment.

[0108] The prepared gradient cemented carbide has a porosity of A04B02C00, a density of 12.83 g / cm 3 , a transverse rupture strength of 2033 MPa, and a hardness (Hv) of 1176 MPa.

[0109] The present application is described in detail above with reference to specific embodiments and / or examples and to the accompanying drawings. However, these descriptions and examples do not limit the present application. Any modifications and / or improvements made to the technical solutions of the present application and the embodiments thereof, which do not depart from the spirit and scope of the present application, are deemed to fall within the scope of the present application. The scope of protection of the present application is defined by the appended claims.

Claims

1. A gradient cemented carbide comprising, from outside to inside, a surface layer portion, a subsurface layer portion and an inner portion, the metal elements of the surface layer portion comprising W, Co elements, the metal elements of the subsurface layer portion comprising W, Co, Ti elements, the metal elements of the inner portion comprising W, Co, Ti and Zr elements, the content of Co elements in the gradient cemented carbide being greater than 10wt%, the thickness of the surface layer portion being 60-110μm. The surface layer portion contains no Ti cubic phase and Zr-containing cubic phase, the subsurface layer portion contains Ti cubic phase, and the inner portion contains Ti cubic phase and Zr-containing cubic phase. The method comprises, after grinding Co powder, cubic phase powder, W powder and WC powder, preparing a green body, and then sintering the green body under vacuum and in a protective atmosphere to obtain the gradient cemented carbide.

2. The gradient cemented carbide according to claim 1, characterized in that, The cubic phase powder is obtained by mixing and grinding TiN powder and ZrC powder, pre-sintering under vacuum, and then sintering in a protective atmosphere.

3. A gradient cemented carbide, characterized in that, The method comprises, after grinding Co powder, cubic phase powder, W powder and WC powder, preparing a green body, and then sintering the green body under vacuum and in a protective atmosphere to obtain the gradient cemented carbide.

5. The method according to claim 4, wherein the weight fraction of Co powder is greater than 10, preferably 13-20, the weight fraction of W powder is 0.1-1.0, preferably 0.2-0.8, and the weight fraction of WC powder is 70-95, preferably 75-90.

4. A method of producing a gradient cemented carbide, characterized in that, After grinding, the mixed slurry is filtered through a 300-500 mesh screen. The sintering temperature under vacuum is 450-1450℃, preferably 500-1400℃, more preferably 550-1350℃, and the sintering time is 1.5-4.5h, preferably 2-4h, more preferably 2.5-3.5h.

8. The method according to claim 4, wherein the sintering temperature in a protective atmosphere is 1300-1600℃, preferably 1350-1550℃, more preferably 1400-1500℃, and the sintering time is 40-90min, preferably 45-80min, more preferably 50-70min. The sintering under vacuum is performed at a heating rate of 1-2℃ / min, at 500-600℃ for 1h, at a vacuum degree of 12-18Pa to remove the forming agent, at a heating rate of 3-5℃ / min, at a sintering temperature of 1200-1300℃ for 1h, at a vacuum degree of 8-12Pa to complete the solid phase stage sintering, and at a heating rate of 2-4℃ / min, at a sintering temperature of 1300-1400℃ for 25-40min, at a vacuum degree of 4-6Pa to complete the liquid phase stage sintering. The gradient cemented carbide is used for manufacturing a turning tool.

6. The method of claim 4, wherein, ​ 7. The method of claim 4, wherein, ​ ​ ​ 9. The method of claim 7, wherein, ​ 10. Use of a gradient cemented carbide according to claim 1, characterised in that, ​

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